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High chromium cast irons: destabilized-subcritical secondary

carbide precipitation and its effect on hardness and wear

properties.

María Agustina Guitar1,a, Sebastián Suárez1, Orlando Prat2, Martín Duarte Guigou3,4,

Valentina Gari3, Gastón Pereira4, Frank Mücklich1

1 Department of Materials Science, Saarland University. Campus D3.3, D-66123,

Saarbrücken, Germany.

2 Departamento de Ingeniería de Materiales, Universidad de Concepción, Edmundo

Larenas 270, Concepción, Chile

3 Programa de Ingeniería de Materiales, Facultad de Ingeniería y Tecnologías,

Universidad Católica del Uruguay, Av. 8 de Octubre 2738 - CP 11600, Montevideo,

Uruguay.

4 Laboratorio de Desarrollo de Nuevos Materiales, Tubacero S.A., Cnel. Raíz 949 -

CP12900, Montevideo, Uruguay.

Corresponding author: a.guitar@mx.uni-saarland.de - Tel./Fax: +49 681 302

70521/70502

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ABSTRACT

This work evaluates the effect of a destabilization treatment combined with a sub-critical

diffusion (SCD) and a subsequent quenching (Q) steps on precipitation of secondary

carbides and their influence on the wear properties of HCCI (16%Cr). The destabilization

of the austenite at high temperature leads to a final microstructure composed by eutectic

and secondary carbides, with an M7C3 nature, embedded in a martensitic matrix. An

improved wear resistance was observed in the SCD+Q samples in comparison to the Q

one, which was attributed to the size of secondary carbides.

Keywords: High chromium white cast iron – Solid state transformation – secondary

carbides precipitation – Wear - microstructure

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Introduction

High Chromium Cast Irons (HCCI) have been a material of choice for wear resistant

components in the mining and mineral processing industries for a long time, given their

outstanding wear and erosion resistance [1]. The main reason for such resistance can be

found in the hard MnCm reinforcement carbides embedded within the metallic matrix,

being those of hypoeutectic, eutectic or hypereutectic origin [2,3]. Hypoeutectic

compositions are the main group of materials in use in the industry, with compositions

described in the ASTM standard A532 groups II and III [4]. Although hypereutectic

HCCI would show a higher carbide volume fraction than hypoeutectic HCCI and it could

be considered as a much more efficient wear resistant material; however, their low

castability severely hinder their practical manufacture and use [5].

The microstructure and mechanical properties of HCCI are a direct consequence of the

eutectic carbide content, matrix microstructure and the presence of secondary carbides

embedded within the metallic matrix [6]. Particularly, the presence of secondary carbides

has shown to improve the wear resistance behavior of the entire composite [2]. Secondary

and also eutectic carbides characteristics are strongly influenced by the content of its main

constituents (C ,Cr) [2,3,7], minor alloying elements (Mo, Ti, V, Nb, W) [8–12] and

thermal processing [5,13].

The wear resistance and mechanical properties, especially the hardness, of HCCI depend

on the type, morphology and distribution of carbides, and on the nature of the supporting

matrix structure which, in turn, depends on the chemical composition and on any

subsequent thermal treatments [14–16]. The secondary carbides, may increase the wear

resistance of the matrix itself, providing load support and reducing damage by third

bodies scratching the surface, by the same mechanisms that has been studied carefully

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with more conventional, low reinforcement volume fraction, composite materials [17,18].

To some extent, the precipitation of secondary carbides should lead to a more bimodal–

like reinforced composite behavior, which has been proved previously to have a better

wear performance than its mono modal counterparts [19,20].

Besides, the secondary carbides which precipitate in the matrix regions also influence

wear resistance by increasing the matrix strength through dispersion strengthening and

microstructure refinement, the fine secondary carbides can increase the mechanical

support of the eutectic carbides. [10,11,21].

Precipitation of secondary carbides during a sub-critical destabilization treatment [22] has

been previously described for HCCI and the usual quenching procedure, i.e. holding

temperatures above the critical line (from 0.5 to 4 hours) before cooling in air. In most

cases the authors compared the latter to a series of sub-critical approaches as references

[11,23,24]. Heat treatments are designed to precipitate a fine dispersion of Fe-Cr carbides

within the matrix. Multi-step treatments have been proposed (martensite formation and

post-quench sub-critical annealing for carbide precipitation), with a significant secondary

carbide precipitation, although the intermediate air cooling induced a sharp decrease in

mechanical properties. The decrease in mechanical properties was associated with the

formation of ferrite phases in the sub-critical step; in general sub-critical treatments have

been associated with hardness reduction [25]. Furthermore, the type, size and distribution

of secondary carbides formed during the destabilization depend on the composition and

destabilization temperature. The secondary carbides are efficient for improving the

abrasion resistance [26,27].

The proposed destabilized plus sub-critical diffusion treatment should lead to a sharper

increase in size and a better distribution of secondary carbides within the matrix, without

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affecting other mechanical properties. Assuming that this material behaves tribologically

as a bimodal size reinforced MMC, the post treatment wear resistance should improve,

despite the slight changes in other mechanical properties [23,28].

In this work, contrary to other works found in the literature [25], a destabilization process

is carried out followed directly by a sub-critical diffusion (SCD) step in order to

precipitate and grow secondary carbides. The main objectives of the present work are: (a)

to analyze if increased secondary carbide precipitation is achievable with a high

temperature destabilization and multi-step sub-critical heat treatment, using limited minor

alloying elements and preserving mechanical properties; (b) to compare such results with

a standard above critical destabilization treatment and (c) to analyze the effect of such

increased carbide precipitation on wear resistance.

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Experimental

HCCI samples were manufactured in arc furnace and casted in cubic sand molds.

Chemical composition was determined by Emission Spectroscopy methods using a GNR

Metal Lab 75/80 Optical Emission Spectroscope. The chemical composition of the

studied HCCI is: C (2.43 wt.%) - Si (0.47 wt.%) – Mn (0.76 wt.%) – Cr (15.84 wt.%) -

Ni (0.41 wt.%) – S (0.02 wt.%) – P (0.02 wt.%) – Cu (0.04 wt.%) – Fe (balance).

The thermal treatments were carried out by two methods: (i) destabilization above critical

temperature (1253 K) and quenching and (ii) destabilization followed by a sub-critical

diffusion (SCD) process at 923 K for 12 h. Destabilization temperature was determined

as critical temperatures plus 50 K and SCD temperature was set at 50 K below zero-

transformation-temperature for the corresponding chemical composition for maximum

carbide precipitation. Thus, observations were carried out in four samples: (1) as cast

conditions, (2) SCD, (3) Quenched (Q) and (4) SCD + Q; being the SCD sample an

intermediate state. Schematic representation for each treatment is shown in Figure 1.

The samples were ground with embedded diamond discs (up to grit 1200) and polished

using diamond powder suspensions up to 1 µm mean diameter. The samples were etched

with Villela´s reagent (1 g picric acid + 5 ml HCl 100 ml ethanol) for 30 seconds at room

temperature. Optical microscopy observations were carried out in all samples using a

Leica CTR6000 microscope and images were acquired using a Jenoptik CCD Camera.

SEM characterization was carried out with a FE-SEM Helios Nanolab 600 (FEI company)

working with an acceleration voltage of 10 kV and a 1.4 nA beam current. For a proper

contrast between phases, a high sensitivity solid state backscattered electrons detector

(vCD) was used. Electron Backscattered Diffraction with an acceleration voltage of 20

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kV and 11 nA beam current together with the TSL OIM Data Collection software were

used for the carbide type identification.

Macro-hardness on the Rockwell C scale was measured using a Wilson durometer with a

diamond indenter and 150 Kgf load. Micro hardness was measured with a Leica Vickers

microindenter applying 0.500 Kgf load on all samples for both, matrix and primary

carbides. The indentation time in all cases was 15 s.

X-ray diffraction testing was performed with a PANalyical Empyrean X-ray

diffractometer. The diffractograms were obtained using a symmetrical θ―θ geometry

configuration and a Cu Kα1 radiation (λ = 0.15406 nm). The incident and diffracted

optical geometries were parallel and the diffraction angle (2θ) was varied from 25° to

100° with a step size of 0.013° and a 50 s/step rate. The applied voltage and current were

40 kV and 40 mA, respectively. For the phase identification and indexing, the High Score

Plus software and ICDD Database were used.

Wear test was carried out in a CSM macro tribometer in rotational mode using an alumina

pin, applying a 5N load for 4000 laps with 2.46mm radius at a speed of 5.0 cm/s with

environment conditions set at 298 K (25 °C) and 45% humidity. Wear tracks were

observed using a Zygo NewView 7300 white light interferometer (for wear volume

calculations) and SEM microscopy.

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Results and Discussion

Phase and microstructural evolution

Figure 2 shows the XRD measurements for the different states of the treated samples. The

as-cast condition is represented by the eutectic carbide embedded in an austenitic matrix

(M7C3/Fe-γ), as described in [25] and also predicted in the Jackson´s diagram for the Fe-

Cr-C system [29]. Some authors also reported the presence of martensite surrounding the

eutectic carbide particles as a result of the excessive consumption of C and Cr by the

eutectic carbides [25]. Consequently, the matrix remain depleted in these elements,

resulting in the increase of the Ms temperature and promoting the martensite formation.

The SCD treatment lead to the transformation of the matrix from austenite to ferrite and

carbides of the form M7C3 were detected. XRD measurements confirms the absence of

martensite in this sample, whose presence is not expected since the sample was cooled in

air after the sub-critical diffusion treatment. On other way, the quenching of the material

(Q sample) generates a complete martensitic transformation of the matrix. Very small

peaks corresponding to M7C3 are also observed. The SCD+Q sample showed, just like

the Q sample, a matrix formed by martensite and carbides of the form M7C3. In the inset

of Figure 2, the most intense peaks from ferrite and martensite are shown. A shift in the

diffraction angle and an increasing width of the peak, as a consequence of the lattice

distortion, allow the identification of martensite form ferrite.

In order to corroborate the obtained phases in the Q and SCD+Q samples, thermodynamic

calculations with the Thermo-Calc software and EBSD measurements were performed,

see Figure 3. The simulation was performed using all the elements of the studied alloy

(as shown in Table 1), and using the TCFe8 database. Figure 3a shows the fraction of the

stable phases in the system at different temperatures. From the thermodynamic point of

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view, it is possible to state that the only stable precipitate within the performed heat

treatment temperature range in the alloy is of the type M7C3, validating the results

obtained by XRD. Moreover, EBSD was used for the characterization of individual

secondary carbides (about 40 secondary carbides were randomly measured). In all cases

an hexagonal crystallographic structure was detected, which corresponds to the M7C3

carbides (Figure 3b). The precipitation of M23C6-type carbides might only be possible at

temperatures under 738 K (465°C), whereas cementite (M3C) is not expected to

precipitate due to the high amount of Cr. Furthermore, it is well known that carbide-

forming elements replace the less stable cementite, which dissolves into a finer alloy

carbide dispersion [21]. This replacement by the alloy carbides, which is more resistant

to coarsening, is able to produce an increase in hardness at higher temperatures [21]. The

strengthening efficiency of this carbide dispersion will depend on the refinement of the

precipitates and their volume fraction.

Figure 4 shows the microstructure for the as cast, SCD, Q and SCD+Q samples. The as

cast microstructure is composed by primary austenite matrix surrounded by the eutectic

mixture of M7C3 carbides/austenite, as also indicated in the diffractogram of Figure 2.

Given the Cr and C content, no secondary carbides are expected to be observed in this

sample. This microstructure is the starting point for the subsequent heat treatments

previously described in Figure 1. SCD process produces a microstructure of ferrite and

secondary carbides. The etching employed here (Villela´s reagent) attacks the ferrite-

carbide microstructure, therefore the optical image in Figure 4b shows dark regions

corresponding to this type of microstructure, which is result of the high-concentration of

small secondary carbides. The Q and SCD+Q microstructures (Figure 4c and Figure 4d)

show a fine martensitic matrix surrounding the unchanged eutectic carbide phase when

compared with as cast conditions. Small secondary carbides are expected to be present in

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these samples, however, they cannot be resolved with optical microscopy techniques.

Black and white arrows in Figure 4 indicate the regions of eutectic and secondary

carbides, respectively. For better resolution and details of presence and distribution of

secondary carbides, SEM images were acquired (Figure 5).

SEM images in BSE mode for all the samples taken at different scales are shown in Figure

5. As previously described, no secondary carbides are observed in the austenitic matrix

of the as-cast condition. The destabilization treatment above the critical temperature 1253

K (980 °C) allow the precipitation of secondary carbides, as shown in SCD (Figure 5b),

Q (Figure 5c) and SCD+Q (Figure 5d) samples. As learned from the XRD results, in the

SCD sample, the secondary carbides are embedded in a ferritic matrix, whereas in the Q

and SCD+Q samples in a martensitic matrix. In all cases the secondary carbides have

apparently randomly nucleated and they show a round-like morphology.

The addition of a sub-critical diffusion (SCD) process following the destabilization step

caused visible modifications compared to the as-cast material and also to quenched

material. Large quantities of secondary carbides were precipitated within the austenitic

matrix, which has transformed completely to ferrite by holding the temperature at 923 K

(650 °C) (see Figure 2). Secondary carbides of different sizes can be observed embedded

in the matrix. The prolonged duration of the SCD process allows the precipitation of a

large quantity of new carbides and the growth of those precipitated during the

destabilization step. This is clearly observable comparing Figure 5b and Figure 5c.

The Q and SCD+Q samples do not show a noticeable change in the amount of secondary

carbides. However, the SCD+Q sample (Figure 5d) shows an apparent increase in size of

secondary carbides inside the martensitic matrix (compared to the Q sample). The

secondary carbides size was determined after image analysis (I-A) of SEM (BSE)

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micrographs using the software A4i©. The calculated average carbide size was 0.44 µm

and 0.51 µm for the Q and SCD+Q samples, respectively. The combination of a SCD

state and quenching provides a coarse-sized secondary carbide distribution throughout

the matrix, as indicated in the carbide size distribution of Figure 5 and also supported by

the evidence observed by electron microscopy (Figure 4). This results from the sub-

critical diffusion treatment which allows the secondary carbides to increase in size,

mainly due to an extended time of energy input to the system.

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Hardness and wear response evaluation

It is well known that the microstructure plays a fundamental role in the resulting

mechanical properties. The wear resistance and mechanical properties, especially the

hardness of HCCI depend on the type, morphology and distribution of carbides, and on

the nature of the supporting matrix structure which, in turn, depends on the chemical

composition and subsequent thermal treatments [14–16]. For this reason, the influence of

the microstructure after thermal treatments in the wear response are being analyzed in this

section.

The mechanical properties play a fundamental role in the applicability of these materials

in the field. They would determine their machinability and subsequently, define the

application in components. Specifically, microhardness measurements show the

variations of local mechanical properties which may significantly influence the abrasion

behavior of the material [30]. The macro and microhardness values of three different

stages are represented in Figure 7a and Figure 7b, respectively. The expected change in

macro-hardness due to the austenitic (as-cast sample) to martensitic (Q and SCD+Q

samples) matrix change is clearly noticeable (Figure 7a). However, the presence of

secondary carbides might not be dismissed as an influencing factor for this improvement.

However, when comparing the results for the Q and SCD+Q samples, no significant

changes are observed, since both present the same type of matrix. Thus, it is consequential

to state that the main role regarding the mechanical properties is played by the matrix

transformation. Micro-hardness measurements where performed on both, the matrix and

the eutectic carbides. The same trend as in the macro-hardness measurements was

observed when only the matrix is evaluated, as shown in Figure 7b. Additionally, a slight

but not statistically representative increase in the matrix micro-hardness of the SCD+Q

sample can be observed in comparison to the Q sample. However, when the eutectic

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carbides zone is evaluated, both the carbides’ intrinsic mechanical response and the

surrounding matrix behavior are considered as a whole system due to the indentation size

effects (Figure 7c). The same increase in micro-hardness is observable in the eutectic

carbide areas of the SCD+Q respect to the Q sample. Those similar hardness values in Q

and SCD+Q is mainly due to the precipitation of the same type of secondary carbides.

When comparing the micro-hardness of the eutectic carbide-containing regions of both

heat treated samples to the as cast condition, the difference can be attributed to the

increase in the supporting role of the matrix surrounding the eutectic carbides. It is worth

noting that in all cases, an incoherent interface between the M7C3 phase and the matrix is

expected, since the interatomic distances show a misfit larger than 25% [21]. This would

further account for an efficient particle dispersion strengthening due to the effective

hindering of dislocation mobility. Additionally, it has been reported that the quenching

step adds a thermal strain to the matrix, which results in a dislocation punching during

plastic strain and accounts for a significant strengthening [31].

Regarding the tribological behavior of the materials, the wear rates after pin on disc tests

are shown in Figure 8. A significant reduction in wear rate of the Q and SCD+Q samples

is observed when compared to the as-cast condition, being 45 and 83%, respectively.

These results correlate quite well with the hypothesis that harder materials tend to show

lower wear rates, given by the conventional Archard behavior/relationship [32].

However, both heat treated samples show a strongly differentiated wear behavior.

Specifically, the SCD + Q sample showed a reduction of 69% compared to the quenched

state.

This enhanced wear resistance of the SCD+Q sample cannot be explained by the macro-

and micro-hardness values. It has already been reported that the wear behavior shows a

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weak dependency on the macrostructure [26]. Thus, for the wear response of these type

of materials (which should be considered as composite materials), not only the carbide

structure or distribution should be considered, but also the nature of the matrix

surrounding the carbides and their interaction as a whole must be taken into account. This

synergistic effect is well-known in multimodal systems [19]. An important role in the

improvement in wear response of Q and SCD+Q samples is played, as already shown for

the hardness, by the austenite/martensite transformation of the matrix added to the

secondary carbides precipitated on the treated samples. Furthermore, even in cases where

a full martensitic transformation is not achieved, the amount of retained austenite might

enhance the response by acting in two ways: inactivating the crack tip formation and

hindering the crack development coming from more brittle phases [27].

During the wear test, oxidation occurs predominantly on the matrix parts of the track

(Figure 9), whereas carbides are usually very slightly affected [33]. This oxidation leads

to a continuous generation of hard and brittle wear particles that are consistently

introduced to the sliding contact, thus inducing an oxide-driven transition from a mild to

a severe third-body wear activity. Additionally, it has been reported for related systems

that the wear rate decreases either with increasing carbide volume fraction (CVF) or with

the increasing in carbide size for the same CVF [26]. In both heat treated samples, the

total CVF is expected to be the same, since they present the same chemical composition.

Yet, some difference was observed in the size distribution of the secondary carbides of

both samples (as observed in Figure 6), being larger in the SCD+Q sample. Hence, these

larger carbides provide an effective protection to the matrix (more wear-prone) from the

abrasion. Additionally, it has been observed a higher crack density in carbides with the

larger axis parallel to the wear surface [26], indicating a strong capacity of the carbides

to absorb energy during friction. Then, it would be reasonable to expect that the sample

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that shows a larger mean secondary carbide size will provide larger non-reactive areas of

load support (in this case, the SCD+Q sample), would therefore present a lower wear rate.

Summarizing, this work shows that the application of a multi-step thermal treatment of a

high chromium cast iron consisting of the austenite destabilization and a subsequent

quenching, lead to the precipitation of carbides inducing a significant reduction in the

wear loss by an accurate tailoring of the secondary phase. When compared to a traditional

quenching process, the same matrix phase is obtained but with different mean secondary

carbide size, which further translates into better wear response without sacrificing the

overall mechanical properties.

Concluding remarks

In this study, we evaluated the effect of a destabilization treatment combined with a sub-

critical diffusion and a quenching steps in the secondary carbides precipitation and their

influence in the wear properties. It was observed, that HCCI in the as-cast condition

composed by an austenitic matrix and eutectic carbides are not sufficient for a good

abrasive wear resistance. In all cases, secondary carbides precipitated within the matrix

are needed for an increase in wear response. A destabilization process followed by a SCD

step allow not only the precipitation but also the growth of secondary carbides. An

additional destabilization and quenching stages transform the matrix into martensite,

resulting in a matrix with a higher hardness but no observable changes in secondary

carbides. This process sequence also avoids the presence of ferrite (or pearlite) in the final

microstructure, which would be detrimental for the mechanical properties. Finally, an

improved wear resistance was observed in the SCD+Q sample compared to the Q one,

being the size of secondary carbides apparently responsible for this behavior. For this, the

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destabilization followed by a SCD is a fundamental step for the secondary carbides

growth.

Wear mechanisms will be thoroughly analyzed in a further work, in order to evaluate the

effect of microstructure changes in the wear rate and wear mechanisms, principally those

related to the secondary carbides nature, shape and size.

Acknowledgements

This work was supported by the CREATE-Network Project, Horizon 2020 Program of

the European Commission (RISE Project N° 644013).

The authors want to thank to Priv.-Doz. Dr. Jose Garcia from AB Sandvik Coromant

R&D (Technology Area Manager Carbide and Sintering) for the helpful discussion and

verification of some of the results.

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Met. Sci. Heat Treat. 55 (2013) 181–183.

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List of Figures

Figure 1: schematic representation for the different thermal treatments (a) quenching,

(b) sub-critical diffusion and (c) sub-critical diffusion followed by quenching.

Figure 2: XRD diffractograms for the different thermal treatments. The inset in the image

shows the shift and peak width of the martensite diffraction peaks (samples Q and

SCD+Q) respect to that from the ferrite (sample SCD).

Figure 3: a) Phase fraction vs temperature diagram, calculated with Thermocalc and b)

Kikuchi patterns corresponding to the secondary carbides showing the hexagonal

crystallographic structure (M7C3 type).

Figure 4: Optical microscopy images showing: (a) the as cast condition; (b) the SCD

sample (c) the Q sample; (d) the SCD+Q sample. The samples were etched with Vilella

reagent for matrix/carbide contrast. Black arrows indicate the regions of eutectic

carbides and the white arrows indicate regions of secondary carbides.

Figure 5: SEM (BSE) Images for all the samples at different magnifications from left to

right 500X, 2000X and 6500X, respectively. a) as cast, b) SCD, c) Q and d) SCD+Q.

Figure 6: secondary carbides size distribution obtained after image analysis of SEM

(BSE) images.

Figure 7: (a) Hardness test results in the Rockwell C Scale and (b) micro-hardness

measurements for three microstructural states: as Cast, Q and SCD + Q. Standard

deviation of the measures is also plotted. (c) Optical micrograph showing the area

influenced by the micro indentation during the eutectic zone measurement.

Figure 8: wear rate volume in mm2.N-1 calculated after pin on disc test for the different

treated samples.

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Figure 9: (a) Backscattered electron micrograph of the wear track in the As-Cast sample

(b) wear track of the quenched sample. The darker areas are tribologically induced

oxides, whereas the lighter regions are either carbides or unmodified matrix material.

The white arrows indicate regions where the eutectic carbides remain unworn.

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Figure 1

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Figure 2

Figure 3

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Figure 4

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Figure 5

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Figure 6

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Figure 7

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Figure 8

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Figure 9